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PAPER
Cite this: Phys. Chem. Chem. Phys.,
2017, 19, 828
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Correlation of annealing time with crystal
structure, composition, and electronic properties
of CH3NH3PbI3xClx mixed-halide perovskite films†
Maryline Ralaiarisoa,a Yan Busby,b Johannes Frisch,c Ingo Salzmann,a
Jean-Jacques Pireauxb and Norbert Koch*c
Using 3D imaging with time-of-flight secondary ion mass spectrometry (ToF-SIMS) complemented by
grazing-incidence X-ray diffraction (GIXRD), we spatially resolve changes in both the composition and
structure of CH3NH3I3xClx perovskite films on conducting polymer substrates at different annealing
stages, in particular, before and after complete perovskite crystallization. The early stage of annealing is
characterized by phase separation throughout the entire film into domains with perovskite and domains
with a dominating chloride-rich phase. After sufficiently long annealing, one single perovskite phase of
homogeneous composition on the (lateral) micrometer scale is observed, along with pronounced film
texture. This composition evolution is accompanied by diffusion of chloride from the perovskite layer
Received 14th September 2016,
Accepted 30th November 2016
towards the conducting polymer substrate, and even accumulation there. Photoelectron spectroscopy
DOI: 10.1039/c6cp06347k
full conversion, which correlates with the change of film composition. Our results accentuate the
importance of chloride for the formation of crystalline and textured films, which are crucial for
www.rsc.org/pccp
enhancing the PV performance of perovskite-based solar cells.
analysis further shows that perovskite films become increasingly n-type with annealing time and upon
1. Introduction
For the past six years, organometal halide perovskites have
spurred the field of photovoltaics, the power-conversion efficiency (PCE) of perovskite based solar cells having exceeded
20%1 in 2016. Most interest was attracted by methyl ammonium
(MA) lead halide perovskite CH3NH3PbX3 (MAPbX3; X = I, Cl, or a
mixture thereof), generally obtained from two compounds: the
organic methyl ammonium halide (MAX) and the inorganic lead
halide (PbX2). Crystalline perovskite films are grown therefrom
following a vast range of deposition methods2–5 including, in
particular, a simple one-step solution process.6 Today, their
outstanding performance and applicability, including mechanical
flexibility, make perovskite-based photovoltaic cells a highly
promising future alternative7–9 as well as add-on10 to conventional silicon-based photovoltaics.
a
Institut für Physik & IRIS Adlershof, Humboldt-Universität zu Berlin,
Brook-Taylor-Str. 6, 12489 Berlin, Germany
b
Research Center in the Physics of Matter and Radiation (PMR), Laboratoire
Interdisciplinaire de Spectroscopie Electronique (LISE), University of Namur,
Rue Joseph Grafé 2, 5000 Namur, Belgium
c
Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Albert-Einstein-Str. 15,
12489 Berlin, Germany. E-mail: [email protected]
† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6cp06347k
828 | Phys. Chem. Chem. Phys., 2017, 19, 828--836
The widely used mixed-halide perovskite MAPbI3xClx, (where x
refers to the final chlorine amount), is typically obtained from
methyl ammonium iodide (MAI) and lead chloride (PbCl2) in a
molar ratio of 3 : 1. In contrast to the single halide perovskite
MAPbI3, as obtained from MAI and lead iodide (PbI2),
the incorporation of chloride turned out to induce long-range
film crystallinity,11 significantly improved charge-transport
properties12,13 and charge carriers diffusion lengths beyond
1 mm, which is one order of magnitude longer than that in
MAPbI3.14 Solution preparation method from pre-formed
MAPbI3xClx powder has recently led to highly uniform and
reproducible films with improved efficiency.15 However,
although mixed-halide perovskite films are obtained from a
chlorine-based precursor, the final films consist predominantly
of MAPbI3.9,11,16 The fact that chlorine was not (or only hardly)
detectable in the final perovskite film is repeatedly referred to
as ‘‘chlorine loss’’.8,12,17,18 In this context, chlorine has been
shown to tend to leave the film through the surface by sublimation of MACl during the thermal annealing process, thereby
facilitating the release of the excess of MA.13,19,20 However, the
presence of chlorine inside the film bulk has been frequently
suggested, and was supposed to play a key role for the electronic
properties at the perovskite/substrate interface.21,22 Several investigations addressing various details of the annealing process have
been carried out to date,13,23–25 and they all converge to the
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importance of the formation of an intermediate phase during
the crystallization process. Moreover, the crystallization of the
mixed halide perovskite films has been reported to depend not
only on the annealing temperature but also on the temperature
of the precursor solution before deposition.26 The critical
impact of chlorine is, however, still not fully understood.
Based on a vivid scientific discourse, a model has been
proposed to describe the impact of chloride across the breadth
of perovskite films,17 focusing on its effect on the crystal growth
process.27 Notably, only very few studies have been directed
towards depth profiling the perovskite bulk properties, in
particular, regarding potential changes in the composition
during the process of film formation and annealing
processes.28,29 Such an investigation is, however, essential
both for fully rationalizing the behavior and role of chloride
during the mixed-halide perovskite formation and for characterizing more in detail the mechanisms governing the crystallization in perovskite films.
In the present work, we investigate the three dimensional
(3D) composition of mixed-halide perovskite thin films by nondestructive profiling of the individual components by laterallyresolved time-of-flight secondary ion mass spectrometry (ToF-SIMS)
during, and after the complete film annealing procedure. ToF-SIMS
results are correlated to the crystal structure and morphology
evolution, as obtained from grazing-incidence X-ray diffraction
(GIXRD) and atomic force microscopy (AFM) analysis. Finally,
annealing-induced changes in the composition are correlated
with the evolution of the electronic structure, as determined by
X-ray and ultraviolet photoelectron spectroscopy (XPS, UPS).
We unequivocally uncover that prior to the completion of
perovskite formation (characterized by the sole presence of
the CH3NH3PbI3 phase) the annealing process is associated
with the formation of two kinds of spatially well-separated
domains throughout the film depth, one corresponding to
the CH3NH3PbI3 phase and another to a lead chloride-rich
phase. Furthermore, in the final perovskite film, chlorine is
found to diffuse mostly towards its interface to the prototypical
underlying hole transport layer [the conductive polymer poly(3,4ethylenedioxythiophene)/poly(styrenesulfonate) – PEDOT:PSS]
with the highest chlorine concentration found actually within
the PEDOT:PSS layer. The accurate determination of the chlorine
location and distribution in the perovskite films also corroborates its role not only during the mixed-halide perovskite
film formation process but also in the final film. Overall, the
structure, the composition, and the electronic properties at the
annealing times investigated are significantly different, thus,
pointing out the key importance of a proper annealing duration
for the formation of mixed-halide perovskite films.
2. Experimental methods
2.1.
Sample preparation
Methyl ammonium (MA) lead iodide chloride CH3NH3PbI3xClx
(MAPbI3xClx) precursor solutions of 40 wt% were obtained
by dissolving CH3NH3I and PbCl2 in a 3 : 1 molar ratio in
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N 0 N-dimethylformamide (DMF). The solution was stirred at
60 1C overnight in a nitrogen glove box.
Indium-tin oxide (ITO) coated glass substrates were solventcleaned and UV-treated for 40 min before spin-coating PEDOT:PSS
(AI4083) at 40 rounds per second (rps) for 30 s. The PEDOT:PSS
films were then annealed at 180 1C for 10 min under ambient
conditions.
The MAPbI3xClx samples were prepared each by spin-coating
(77 rps, 45 s) 25 mL MAPbI3xClx precursor solution onto a 10 mm
10 mm PEDOT:PSS/ITO substrate and subsequently annealed at
100 1C in a nitrogen glove box for 10 min (samples termed S-10)
up to 100 min (termed S-100), respectively. The thickness of the
films obtained is approximately 250 nm. Separate sets of samples
were prepared following the same procedure for each experimental analysis (ToF-SIMS, GIXRD, AFM, XPS and UPS).
2.2.
ToF-SIMS analysis
ToF-SIMS 3D analysis was performed with a dual beam Iontof-IV
instrument in non-interlaced mode by using a Bi3+ primary beam
(25 keV) for the analysis step followed by the sputtering step
(during 10 s and 5 s for S-10 and S-100, respectively) with a lowenergy (500 eV) Cs+ beam. The depth profile is built by alternating
sputtering and acquisition steps on typically 250 250 mm2 and
125 125 mm2 analysis areas for S-10 and S-100, respectively.
At each acquisition, the total secondary ions spectrum (1 o m/z o
800 a.u.) is collected from the few topmost atomic layers. Once
the molecular species are identified, semi-quantitative ToF-SIMS
profiles are reconstructed showing the selected ions’ intensities as
a function of sputter time (depth). The spatial (XY) distribution of
an ion is also reconstructed in 2D color-scaled XY spectral maps,
displaying the depth-integrated intensity of the respective ion with
about 1 mm lateral resolution. Cross-section analysis was done
by selecting X or Y cuts from the XY map, and visualizing the
corresponding XZ or YZ ion depth distribution. The depth
profiling process with the corresponding reconstruction is
illustrated in a simplified way in Fig. S9 of the ESI.†
2.3.
XPS and UPS analysis
Photoemission measurements were performed using an Al Ka
X-ray source (excitation energy: 1486.6 eV) for XPS and a
He-discharge UV source (Omicron) with an excitation energy
of 21.2 eV for UPS, spectra recorded with a Phoibos 150 (Specs)
hemispherical analyzer. The base pressure in the analysis chamber
was o109 mbar. Samples for PES were directly transferred from a
N2-filled glove box into the analysis chamber without exposure to
air. For valence region spectra, a pass energy of 5 eV was used and
the secondary electron cut-off (SECO) was recorded with 10 V
sample bias (to clear the analyzer work function for work function
determination) at a pass energy of 2 eV. XPS peak fitting was done
employing a mixed Gaussian/Lorentzian peak shape and a Shirleytype background using the software package WINSPEC (University
of Namur).
2.4.
GIXRD measurements
GIXRD experiments were performed at BESSY II of the
Helmholtz-Zentrum Berlin für Materialien und Energie GmbH
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(HZB, Germany) at beamline KMC-2 using a primary beam
wavelength of 0.1 nm and a Vantec 2000 area detector (sample
to detector distance: 336 mm) under N2 flux to reduce beam
damage; the setup allowed covering a range in reciprocal space
of ca. 2 Å1 both in q|| and qz direction in a single experiment.
The integration time was 900 s in all cases.
3. Results and discussion
3.1.
Crystal structure analysis
The evolution of the crystal structure of one-step solutionprocessed CH3NH3PbI3xClx perovskite films was monitored
as a function of increasing annealing time at constant annealing
temperature of 100 1C by GIXRD (Fig. 1).
We investigated pristine (not-annealed) (Fig. 1a) films and
those annealed at 100 1C for 10 min (S-10, Fig. 1b), 30 min
(Fig. 1c), and 100 min (S-100, Fig. 1d). In all cases, GIXRD shows
the growth of the perovskite in pronounced fiber texture, i.e.,
a common plane of its crystal structure is preferentially
oriented parallel to the substrate (with random azimuthal
orientation of the crystallites). This growth mode is deduced
from the presence of regions of strongly enhanced intensity on
the diffraction rings (Debye–Scherrer rings) in Fig. 1a–d, which,
within the error margin of the present experiment, can be
well indexed both with a (pseudo-)cubic and a tetragonal
polymorph, as repeatedly observed for perovskite films.30
Fig. 1e shows the expected peak positions calculated for the
tetragonal polymorph in (001) texture as used to indexing our
data. The observation of the (211) reflection at the indicated
position acts as fingerprint for tetragonal instead of cubic
perovskite. For the latter, this reflection would occur at a far
different position.31 The structure factors calculated for the
(211) and (202) reflections (cf. Fig. 1e), and, therefore, the
expected diffraction intensities, are essentially identical in
the tetragonal crystal structure. In the present GIXRD data,
however, the diffraction intensity at the expected (211) position
(dotted circles in Fig. 1a–d) is negligible compared to that of
(202), which demonstrates that in our samples perovskite
predominantly grows in a (pseudo-)cubic polymorph in (001)texture; the corresponding calculated Debye Scherrer rings
depicted in Fig. 1f match the experimental data reasonably
well. Due to the high surface sensitivity of the technique, no
contributions from the underlying ITO substrate are observed;
GIXRD reference data for the bare PEDOT:PSS/ITO substrate
are shown in the ESI,† Fig. S1.
Apart from the perovskite-derived diffraction pattern, for
both the pristine sample and S-10 we observe additional features
(cf. Fig. 1a and b), which, however, cannot be unambiguously
assigned. Both the crystal structures of pristine CH3NH3I and
PbCl2 cannot be held to index these additional reflections.32,33
In literature,24,34 the additional features were assigned to a
Fig. 1 2D GIXRD patterns of CH3NH3PbI3xClx perovskite films, as obtained from a precursor solution of CH3NH3I and PbCl2 (3 : 1) in DMF, (a) without
annealing, and annealed at 100 1C for (b) 10 min, (c) 30 min, and (d) 100 min. (e) Simulated GIXRD pattern of tetragonal perovskite in (001) texture, and
(f) calculated Debye–Scherrer rings (non-textured case) for (pseudo-)cubic perovskite.28 Arrows in (a and b) exemplarily mark additional diffraction
features assigned to perovskite precursors, dotted circles highlight the expected position of the tetragonal (211) reflection, and the arrow in (d) marks a
ring assigned to PbI2 related to the degraded perovskite.
830 | Phys. Chem. Chem. Phys., 2017, 19, 828--836
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‘‘crystalline precursor structure’’ that readily transforms into
perovskite upon thermal annealing but could not further
be crystallographically specified. Interestingly, one of these
diffraction features at q = 1.1 Å1 does agree with the calculated
position of MAPbCl3,25 which is an intermediate phase
assumed to be crucial for the completion of film formation in
solution-processed mixed-halide perovskite.24 Importantly, for
annealing times over 30 min all precursor features vanish
(Fig. 1c and d) and the film appears to be fully converted into
MAPbI3 perovskite. For S-100 (100 min annealing), a faint ring
at q = 0.9 Å1 indicates the emergence of PbI2,35 a known
degradation product of MAPbI3 upon annealing.24
Since the mixed-halide perovskite film structure at 10 min
annealing time, i.e. before complete film formation, is significantly different from the film structure after 100 min, when the
crystallization is completed and no further structure evolution
is observed (except for growing PbI2 content), we applied
spatially resolved chemical composition analysis to samples
S-10 and S-100 in lateral and vertical direction.
3.2.
Morphology, interface and composition analysis
ToF-SIMS analysis was conducted to characterize the perovskite
sample composition in 3D before (S-10) and after complete
crystallization (S-100). In ToF-SIMS profiles in the surface
normal direction, for each sputter step (corresponding to a
sputter time), the intensity of selected secondary ions over the
analyzed area is represented as a data point in the profile. The
sputter time can be directly translated into the probed sample
depth if the sputtering rate is constant, i.e., within a given
homogeneous material. A low-energy Cs+ beam has been
demonstrated to allow for convenient sputtering rates both
on organic and inorganic materials while preserving a high
yield for characteristic high-mass molecular fragments.36,37
Fig. 2 shows the ToF-SIMS profiles obtained for selected masses
on S-10 and S-100. Fragments were carefully selected to identify
the dominating MAPbI3 perovskite phase (PbI3), the PEDOT:PSS
layer (C4H4S), and the ITO substrate (In2O2) while avoiding
mass overlaps. In the not-fully-crystallized perovskite (S-10), fragment ions from the inorganic compound (PbCl2 or its derivative)
are still present. The distribution of PbCl2 and other related
compounds was assessed by selecting and reconstructing the
PbCl3 signal distribution. CN was chosen to identify the
organic MA; this low-mass fragment, however, is not specific
to a structure as it could either come from the perovskite, or
from other compounds such as MACl, MAI, MAPbCl3, etc.
In the ToF-SIMS profile of S-10 (Fig. 2a), PbI3, PbCl3, and
CN are detected with relatively constant intensities during the
whole sputtering process, in contrast to C4H4S (indicative of
PEDOT:PSS), which shows a drop of intensity after 250 s sputter
time. Thus, the precursor/perovskite layer in S-10 consists of
PbCl3-based compounds, MA, and already formed MAPbI3
perovskite (as inferred from GIXRD). The detection of C4H4S
already from the beginning of the experiment for both samples
S-10 and S-100 (Fig. 2a and b) demonstrates the high sensitivity
(up to 1 ppm) of ToF-SIMS, as no signal related to PEDOT:PSS
was detected by XPS measurements (cf. S 2p core level spectra in
Fig. S2, ESI†). While the intensities of the secondary ions of
both samples are not quantitatively comparable, the trends of
the depth profiles allow drawing qualitative conclusions. In
the ToF-SIMS profile of S-10 (Fig. 2a), the signal from the
PEDOT:PSS raises after few seconds of sputtering, while it
increases only after 250 s for S-100 (Fig. 2b). This observation
implies that the precursor/perovskite layer does not cover the
PEDOT:PSS completely after 10 min annealing. AFM investigations confirm this hypothesis (Fig. 3). The AFM micrograph
from S-10 evidences the presence of pinholes and cavities
(circles and squares in Fig. 3b). The high surface corrugation
of the precursor/perovskite film (root mean square roughness
RS = 46 nm) explains the nearly simultaneous detection of both
perovskite- and the PEDOT:PSS-related ions in ToF-SIMS
(Fig. 2a). The sputtering process tends to carry on the high
roughness of the initial surface. In contrast, the AFM micrograph of a not-annealed sample (Fig. 3a) shows needle structures
with large voids in between; this translates into a substantially
higher roughness (RS = 90 nm). With increasing annealing time
the roughness of the films is significantly reduced (S-10: 46 nm,
S-100: 23 nm), as shown in Fig. 3b and c. Annealing the sample
leads to a more homogeneous surface coverage, thus explaining
Fig. 2 ToF-SIMS depth profiles of sample (a) S-10 and (b) S-100: interfaces between the different layers are identified by looking at the intensities of
characteristic ion (PbI3, C4H4S, and In2O2 for the perovskite, PEDOT:PSS, and ITO layers, respectively), as shown by the vertical lines (intensity Z80%
of its respective maximum value).
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Fig. 3 AFM micrographs of samples annealed during (a) 0 min, (b) 10 min (circles and squares indicate pinholes and cavities, respectively) and (c) 100 min
with an overall root mean square roughness RS of 90 nm (0 min), 46 nm (10 min) and 23 nm (100 min), respectively.
the initially much lower signal from the PEDOT:PSS layer in the
ToF-SIMS profile of S-100, as mentioned above.
In contrast to S-10, we see from the ToF-SIMS profiles in
Fig. 2b that the interfaces between the individual layers of S-100
can be better discriminated (see vertical lines in Fig. 2b), where
each layer is identified by the region where the corresponding
ion (perovskite: PbI3, PEDOT:PSS: C4H4S, ITO: In2O2) has
intensities higher than 80% of its maximum value. The interface
with PEDOT:PSS marked by the C4H4S intensity plateau is
reached in S-100 after about 300 s sputtering (Fig. 2b) while it
appears almost from the beginning of the depth profile in S-10.
Importantly, while the chlorine concentration in the
perovskite layer deduced from the ToF-SIMS depth profiles
for S-100 (Fig. 2b) is only slightly above the detection limit
and can be well ascribed to contamination, we find a significant
accumulation of chlorine (37Cl signal) from the precursor/
perovskite within the PEDOT:PSS layer (C4H4S signal) and
at the interface, as both signals have comparable trends.
As an aside we note that it was argued before whether the term
‘‘mixed halide perovskite’’ is still tenable for perovskite
obtained from chlorine-based precursors.12,38 Despite our
results for the S-100 film, where chlorine is found to be only
present at the substrate interface, we retain this wording here,
in order to highlight the relevance of chlorine in the formation
process of perovskite. Since the Cl ionization yield is very high
and the ToF-SIMS technique is semi-quantitative (as applied
here), it is worthy to stress that in a reference film of pristine
PEDOT:PSS (i.e., without a precursor/perovskite overlayer) on
ITO, an 18 times lower relative chlorine content is found (ESI,†
Fig. S3). This excludes that the chlorine signal measured in the
PEDOT:PSS layer of S-100 is due to contamination. For S-10, the
ToF-SIMS profile of 37Cl follows that of PbCl3 suggesting that
both ions represent the same species (ESI,† Fig. S4). For S-100,
the profile of CN has the same trend as that of 37Cl (Fig. 2b),
thus indicating that chlorine is present most likely in the form
of MACl both at the precursor/perovskite-PEDOT:PSS interface
and within the PEDOT:PSS layer.
It has been reported that even a low chlorine concentration
allegedly can significantly improve the charge-transport properties
of perovskite films.12 Particularly, the presence of chlorine at
the perovskite interface to the substrate has been suggested to
promote the charge collection at the interface.21 The ‘‘chlorine
loss’’ in the final mixed-halide MAPbI3xClx perovskite film has
832 | Phys. Chem. Chem. Phys., 2017, 19, 828--836
already been discussed in several studies,13,17,39 where a sublimation of, e.g., MACl, during the annealing step has been
proposed to explain this loss.13,20,39 Fully in line with
literature,13,20 our XPS measurements on S-100 did not reveal
any presence of chlorine above the detection limit (see below).
Nevertheless, the presence of residual chlorine within the final
MAPbI3xClx perovskite has been experimentally assessed by
X-ray absorption near-edge structure (XANES).40 The chlorine
location, however, could not be determined with precision yet.
A strong spatial inhomogeneity in the chlorine concentration
throughout a 60 nm thin perovskite film was deduced from
hard X-ray photoelectron spectroscopy combined with fluorescence yield X-ray absorption spectroscopy.22 To test the theoretical assumption assessing the interfacial role of chloride,
Colella et al. conducted angle-resolved XPS measurement on
a 5 nm thin perovskite film and detected chloride at the
perovskite TiO2 interface.21 However, such a film thickness
does not reflect real device applications. As reported here, the
ToF-SIMS analysis, in contrast to surface sensitive XPS techniques,
allows to track the chlorine content throughout the entire perovskite
layer depth and into the underlying hole transport layer in a sample
with a perovskite film of device-relevant thickness (B250 nm).8,14
This could support the notion of improved charge transport due to
chlorine presence in devices based on MAPbI3xClx perovskite films.
The increase in 37Cl signal intensity down towards and within the
PEDOT:PSS layer itself, as observed here, resembles the chlorine
diffusion observed in mesoporous TiO2 of a perovskite-based solar
cell with the structure (Au/hole-transport-material/perovskite/
TiO2).28 This congruence points towards similar chlorine diffusion
dynamics, independent of substrate and device configuration, and
thus also corroborates the idea that chlorine affects the interfacial
electronic structure by modifying the perovskite/substrate
interface.21 This provides support for the prominence of
mixed-halide MAPbI3xClx perovskite-based solar cells in comparison to their single halide MAPbI3 perovskite counterpart.
Further spatial analysis of the sample composition was
performed by depth-integrating the intensity of the selected
ion for a given area. These 2D ToF-SIMS maps for PbI3, PbCl3
and CN of S-10 (Fig. 4) disclose a lateral phase separation
occurring in the precursor/perovskite layer after 10 min annealing.
Two regions, A and B, can be distinguished: region A (red squares)
is characterized by high PbI3 and low PbCl3 intensities while
the opposite applies in region B (blue circles). Hence, we assign
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Fig. 4 ToF-SIMS 2D (XY) maps taken on S-10 displaying the depth-integrated intensity of the selected ions. The lateral resolution is about 1 mm.
Red squares indicate a region (A in the text) characterized by high PbI3 and CN intensities and a low PbCl3 intensity; blue circles indicate regions
(B in the text) where PbCl3 and CN intensities are high and that of PbI3 low.
region A to the already formed perovskite MAPbI3, the presence of
which was already established by GIXRD analysis (vide supra).
We assume that the high PbCl3 ion intensity in region B
could either be due to the fragmentation of MAPbCl3 or to the
presence of the original inorganic compound PbCl2. Note that
the latter rationale agrees well with the presence of PbCl2
nanoparticles in PbCl2-based perovskite precursor solution, as
reported by Tidhar et al.27 It has been suggested that due to the
lower solubility of PbCl2, such nanoparticles are still present in
the precursor solution and serve as nucleation centers, promoting
the formation of crystalline perovskite.
At this point, results do not allow ruling out the hypothesis
that the measured PbCl3 signal comes from the fragmentation
of MAPbCl3; the presence of MAPbCl3 acting as an intermediate
phase has been indeed confirmed by the GIXRD analysis. To
address this question, the distribution of organic MA halide
species (MAX) in S-10 was monitored by following the high-yield
CN ion XY-distribution (see right panel in Fig. 4). Fig. 4 shows
that regions exist (circles), where PbCl3 and CN signals are
simultaneously high, and, likewise, regions (squares) where this is
true for the PbI2 and CN signals. Therefore, CN is present in
both regions (A, B). In region A, CN may originate from the
fragmentation of MAPbI3 perovskite or MAI (compound in
excess). This latter assumption is further supported by the
XY-reconstruction of CH3NH2 (de-protonated MAI) in Fig. S5
(ESI†) that indicates a spatial distribution very similar to the
one of PbI3. In region B, CN stems from the fragmentation
of methyl ammonium from intermediate compounds such as
MACl or MAPbCl3 related to the crystalline precursor
structure,12,17,24 as observed by GIXRD analysis. Since such an
XY-reconstruction of CH3NH2Cl (MACl compound) could not
be performed due to its extremely low intensity, we assume that
region B is dominated by MAPbCl3 and possibly PbCl2. Our
results are in line with the topotactic transformation proposed
by Moore et al.19 and Williams et al.17 In their model, the
authors suggest the co-presence of a chloride-rich phase with a
peripheral iodide-rich phase during annealing. By diffusion of
MA+ and I from the latter to the former (with the chloride
being volatile), a more stabilized iodide-rich phase evolves with
an enhanced film crystallization, in contrast to chlorine-free
single MAPbI3 perovskite. The intermediate chloride-rich phase
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therefore acts as a preceding structural template for the final
MAPbI3 perovskite. In contrast to the study of Williams et al.17
describing a variation in the vertical composition of mixedhalide perovskite films during annealing, our results point
towards the coexistence of lead chloride-rich and lead iodiderich phases right throughout the film depth (see also the YZ
cross-section map in Fig. S6 in the ESI†), as we schematically
illustrate in Fig. 5a.
These findings underline the importance of chloride for the
formation of highly textured films, which are crucial for enhancing the photovoltaic performance of perovskite-based solar
cells.41 The annealing is associated with the chlorine migration
towards the interface of perovskite and the PEDOT:PSS layer
(which acts as hole transport layer here), while the remaining
Pb recombines with MA+ and I to form the MAPbI3 perovskite. As
shown by the XY maps of S-100 (Fig. S7, ESI†), when the
crystallization is completed, the perovskite film does no longer
show any phase separation (see Fig. 5b).
3.3. Electronic structure and composition analysis by
photoemission
By combining XPS and UPS, we independently explored the
influence of the annealing time not only on the chemical
Fig. 5 Simplified schematic cross-sections as resulted from ToF-SIMS 2D
imaging: (a) illustrates the lateral phase separation between region A,
corresponding to already formed MAPbI3 perovskite and region B, where
perovskite has not formed yet for S-10; and (b) shows a homogeneous
composition and a lower surface roughness for S-100.
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Fig. 6 Photoemission spectra of perovskite samples annealed at 100 1C for 10 min (S-10) and 100 min (S-100). XPS core level spectra of (a) Cl 2p, (b) I 3d,
(c) Pb 4f, and (d) N 1s. UPS (e) SECO spectra and (f) valence spectra.
composition but also on the electronic structure of the
resulting films.
For both S-10 and S-100 samples, no chlorine was detected
from the Cl 2p core level spectra in Fig. 6a, evidencing the
inadequacy of XPS in detecting minute amounts of an element
as opposed to ToF-SIMS (vide supra); further core levels (I 3d,
Pb 4f and N 1s) are shown in Fig. 6b–d. Notably, for S-10
the binding energy (BE) of the I 3d7/2 (at 619.1 eV) and the N 1s
(at 401.9 eV) core levels increased by 0.7 eV, and that of Pb 4f
(at 138.1 eV) by 0.6 eV with prolonged annealing time (S-100).
Very much alike, the valence band onset (VBonset) shifted to the
same direction relative to the Fermi level (EF) by 0.5 eV when
going from S-10 to S-100, and the work function F decreased by
ca. 0.4 eV (Fig. 6e and f). Since all electronic levels (core and
valence) shift towards higher BE from S-10 to S-100 by approximately the same amount, the cause for the shift is most likely
of electrostatic nature. Given the direction of the shift, the
material becomes more n-type upon full conversion to the
perovskite. The small difference in shift between core and valence
levels is accounted for by the concomitant overall change of the
valence electronic structure. The fact that the work function
change is lower than the shift of all electronic levels points
towards a change in the surface electrostatic potential due to
structural/compositional changes in addition to the modified
doping level.
Note that such an electronic-structure modification with
annealing time has been suggested before,25 based on the
argument that an MAI deficit n-dopes perovskite films,42 i.e.,
MA sublimation upon annealing renders the film more n-type.
We further observe a concomitant (small) increase (+0.12 eV) in
the ionization energy (IE = F + VBonset) with annealing time,
which can be related to the variation of the perovskite-film
stoichiometry (in particular, MA+ deficiency) as was recently
suggested.43 This assumption is corroborated by ToF-SIMS
analysis by displaying the CN to PbI3-ratio (CN/PbI3) XY
834 | Phys. Chem. Chem. Phys., 2017, 19, 828--836
image for S-10 and S-100 (Fig. S8, ESI†). In comparison to S-100,
regions exist in S-10 where the relative CN content is higher,
indicating that an overall lower MA content is found after
longer annealing time.
4. Conclusion
The present study provides insights into the nucleation of
the MAPbI3xClx mixed-halide perovskite. The evolution of
structure, morphology, chemical composition, and electronic
structure of precursor and perovskite films was investigated
before, during, and after complete conversion. With annealing
time, solution-processed perovskite films were found to evolve
from an inhomogeneous composition characterized by laterally
separated MAPbI3 perovskite and lead chloride-rich domains,
to a single phase corresponding to a fully crystallized perovskite
after 100 min of annealing at 100 1C. In particular, this very
evolution of the film composition was associated with evident
changes of its crystal structure, which corroborates the crucial
role of PbCl2 for initiating perovskite nucleation and the
importance of intermediate phases, particularly MAPbCl3, for
the formation of highly uniform and textured perovskite film,
which is essential for attaining high-performance perovskite
solar cells.
Relative deficiencies in MA content were found to modify
the electronic structure and promote the formation of more
n-type films with increasing annealing time. Noteworthy, some
chlorine always stays in the annealed films even after the full
crystallization process; this was enabled by the use of ToF-SIMS
with much higher sensitivity than XPS. Before the complete
perovskite conversion, chlorine from the inorganic lead chloride
compound and intermediate phases was detected in isolated
regions as evidenced by ToF-SIMS 2D maps, whereas after
completion of the perovskite film formation, chlorine was mostly
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located at the interface to and within the underlying PEDOT:PSS
layer. These results corroborate not only the key role of chlorine
in the formation of solution-processed mixed-halide MAPbI3xClx
perovskite films but also its suggested implication in the optoelectronic properties of perovskite-based solar cells, thus highlighting the importance of controlling the annealing conditions
and duration for optimizing perovskite film preparation.
Acknowledgements
This work was supported by the Helmholtz Energy Alliance
‘‘Hybrid Photovoltaics’’, the SFB951 (DFG), and the HZB-HU
Graduate School ‘‘hybrid4energy’’.
References
1 M. A. Green, K. Emery, Y. Hishikawa, W. Warta and
E. D. Dunlop, Prog. Photovolt. Res. Appl., 2016, 24, 905–913.
2 J. Burschka, N. Pellet, S.-J. Moon, R. Humphry-Baker,
P. Gao, M. K. Nazeeruddin and M. Grätzel, Nature, 2013,
499, 316–319.
3 M. Liu, M. B. Johnston and H. J. Snaith, Nature, 2013, 501,
395–398.
4 Q. Chen, H. Zhou, Z. Hong, S. Luo, H.-S. Duan, H.-H. Wang,
Y. Liu, G. Li and Y. Yang, J. Am. Chem. Soc., 2014, 136,
622–625.
5 L. Liu, J. A. McLeod, R. Wang, P. Shen and S. Duhm, Appl.
Phys. Lett., 2015, 107, 61904.
6 A. Kojima, K. Teshima, Y. Shirai and T. Miyasaka, J. Am.
Chem. Soc., 2009, 131, 6050–6051.
7 H. J. Snaith, J. Phys. Chem. Lett., 2013, 4, 3623–3630.
8 J. You, Z. Hong, Y. M. Yang, Q. Chen, M. Cai, T.-B. Song,
C.-C. Chen, S. Lu, Y. Liu, H. Zhou and Y. Yang, ACS Nano,
2014, 8, 1674–1680.
9 H. Zhou, Q. Chen, G. Li, S. Luo, T. Song, H.-S. Duan,
Z. Hong, J. You, Y. Liu and Y. Yang, Science, 2014, 345,
542–546.
10 D. P. McMeekin, G. Sadoughi, W. Rehman, G. E. Eperon,
M. Saliba, M. T. Hörantner, A. Haghighirad, N. Sakai,
L. Korte, B. Rech, M. B. Johnston, L. M. Herz and
H. J. Snaith, Science, 2016, 351, 151–155.
11 M. M. Lee, J. Teuscher, T. Miyasaka, T. N. Murakami and
H. J. Snaith, Science, 2012, 338, 643–647.
12 S. Colella, E. Mosconi, P. Fedeli, A. Listorti, F. Gazza,
F. Orlandi, P. Ferro, T. Besagni, A. Rizzo, G. Calestani,
G. Gigli, F. De Angelis and R. Mosca, Chem. Mater., 2013,
25, 4613–4618.
13 H. Yu, F. Wang, F. Xie, W. Li, J. Chen and N. Zhao, Adv.
Funct. Mater., 2014, 24, 7102–7108.
14 S. D. Stranks, G. E. Eperon, G. Grancini, C. Menelaou,
M. J. P. Alcocer, T. Leijtens, L. M. Herz, A. Petrozza and
H. J. Snaith, Science, 2013, 342, 341–344.
15 D. Song, J. H. Heo, H. J. Han, M. S. You and S. H. Im,
J. Power Sources, 2016, 310, 130–136.
This journal is © the Owner Societies 2017
PCCP
16 B. Park, B. Philippe, T. Gustafsson, K. Sveinbjörnsson,
A. Hagfeldt, E. M. J. Johansson and G. Boschloo, Chem.
Mater., 2014, 26, 4466–4471.
17 S. T. Williams, F. Zuo, C.-C. Chueh, C.-Y. Liao, P.-W. Liang
and A. K.-Y. Jen, ACS Nano, 2014, 8, 10640–10654.
18 B. Philippe, B.-W. Park, R. Lindblad, J. Oscarsson,
S. Ahmadi, E. M. J. Johansson and H. Rensmo, Chem.
Mater., 2015, 27, 1720–1731.
19 D. T. Moore, H. Sai, K. W. Tan, L. A. Estroff and U. Wiesner,
APL Mater., 2014, 2, 81802.
20 E. L. Unger, A. R. Bowring, C. J. Tassone, V. L. Pool,
A. Gold-Parker, R. Cheacharoen, K. H. Stone, E. T. Hoke,
M. F. Toney and M. D. McGehee, Chem. Mater., 2014, 26,
7158–7165.
21 S. Colella, E. Mosconi, G. Pellegrino, A. Alberti, V. L. P.
Guerra, S. Masi, A. Listorti, A. Rizzo, G. G. Condorelli,
F. De Angelis and G. Gigli, J. Phys. Chem. Lett., 2014, 5,
3532–3538.
22 D. E. Starr, G. Sadoughi, E. Handick, R. G. Wilks, J. H.
Alsmeier, L. Köhler, M. Gorgoi, H. J. Snaith and M. Bär,
Energy Environ. Sci., 2015, 8, 1609–1615.
23 M. Saliba, K. W. Tan, H. Sai, D. T. Moore, T. Scott, W. Zhang,
L. A. Estroff, U. Wiesner and H. J. Snaith, J. Phys. Chem. C,
2014, 118, 17171–17177.
24 K. W. Tan, D. T. Moore, M. Saliba, H. Sai, L. A. Estroff,
T. Hanrath, H. J. Snaith and U. Wiesner, ACS Nano, 2014, 8,
4730–4739.
25 P. P. Atittaya Naikaew, Appl. Phys. Lett., 2015, 106, 232104.
26 D. Khatiwada, S. Venkatesan, N. Adhikari, A. Dubey, A. F.
Mitul, L. Mohammad, A. Iefanova, S. B. Darling and Q. Qiao,
J. Phys. Chem. C, 2015, 119, 25747–25753.
27 Y. Tidhar, E. Edri, H. Weissman, D. Zohar, G. Hodes,
D. Cahen, B. Rybtchinski and S. Kirmayer, J. Am. Chem.
Soc., 2014, 136, 13249–13256.
28 F. Matteocci, Y. Busby, J.-J. Pireaux, G. Divitini, S. Cacovich,
C. Ducati and A. Di Carlo, ACS Appl. Mater. Interfaces, 2015,
7, 26176–26183.
29 B. Yang, J. Keum, O. S. Ovchinnikova, A. Belianinov,
S. Chen, M.-H. Du, I. N. Ivanov, C. M. Rouleau, D. B.
Geohegan and K. Xiao, J. Am. Chem. Soc., 2016, 138,
5028–5035.
30 C. C. Stoumpos, C. D. Malliakas and M. G. Kanatzidis,
Inorg. Chem., 2013, 52, 9019–9038.
31 T. Baikie, Y. Fang, J. M. Kadro, M. Schreyer, F. Wei, S. G.
Mhaisalkar, M. Graetzel and T. J. White, J. Mater. Chem. A,
2013, 1, 5628–5641.
32 O. Yamamuro, T. Matsuo, H. Suga, W. I. F. David,
R. M. Ibberson and A. J. Leadbetter, Acta Crystallogr., Sect.
B: Struct. Sci., 1992, 48, 329–336.
33 R. L. Sass, E. B. Brackett and T. E. Brackett, J. Phys. Chem.,
1963, 67, 2863–2864.
34 Y.-C. Huang, C.-S. Tsao, Y.-J. Cho, K.-C. Chen, K.-M. Chiang,
S.-Y. Hsiao, C.-W. Chen, C.-J. Su, U.-S. Jeng and H.-W. Lin,
Sci. Rep., 2015, 5, 13657.
35 Z. Zheng, A. Liu, S. Wang, Y. Wang, Z. Li, W. M. Lau and
L. Zhang, J. Mater. Chem., 2005, 15, 4555–4559.
Phys. Chem. Chem. Phys., 2017, 19, 828--836 | 835
View Article Online
PCCP
40 V. L. Pool, A. Gold-Parker, M. D. McGehee and M. F. Toney,
Chem. Mater., 2015, 27, 7240–7243.
41 P. Docampo, F. C. Hanusch, N. Giesbrecht, P. Angloher,
A. Ivanova and T. Bein, APL Mater., 2014, 2, 81508.
42 Q. Wang, Y. Shao, H. Xie, L. Lyu, X. Liu, Y. Gao and
J. Huang, Appl. Phys. Lett., 2014, 105, 163508.
43 J. Emara, T. Schnier, N. Pourdavoud, T. Riedl, K. Meerholz
and S. Olthof, Adv. Mater., 2016, 28, 553–559.
Open Access Article. Published on 30 November 2016. Downloaded on 18/06/2017 05:27:33.
This article is licensed under a Creative Commons Attribution 3.0 Unported Licence.
36 Y. Busby, S. Nau, S. Sax, E. J. W. List-Kratochvil, J. Novak,
R. Banerjee, F. Schreiber and J.-J. Pireaux, J. Appl. Phys.,
2015, 118, 75501.
37 C. Noël and L. Houssiau, J. Am. Soc. Mass Spectrom., 2016,
27, 908–916.
38 M. Grätzel, Nat. Mater., 2014, 13, 838–842.
39 A. Dualeh, N. Tétreault, T. Moehl, P. Gao, M. K. Nazeeruddin
and M. Grätzel, Adv. Funct. Mater., 2014, 24, 3250–3258.
Paper
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